Tuning the Reversibility of Oxygen Redox in Lithium-Rich Layered

Feb 28, 2017 - Recently, more and more new high-capacity lithium-rich layered oxides involving both metal and oxygen redox have been proposed. However...
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Tuning the Reversibility of Oxygen Redox in Lithium-Rich Layered Oxides Biao Li, Huijun Yan, Yuxuan Zuo, and Dingguo Xia* Beijing Key Laboratory of Theory and Technology for Advanced Batteries Materials, College of Engineering, Peking University, Beijing 100871, P. R. China S Supporting Information *

ABSTRACT: Recently, more and more new high-capacity lithium-rich layered oxides involving both metal and oxygen redox have been proposed. However, the structural stability was influenced by the irreversible oxygen redox, which leads to the instability of the oxygen framework. Here, we propose that the reversibility of the oxygen redox in Li2RuO3 can be controlled by tuning its electronic structure via incorporating boron atoms into the interstitial sites of the Li2RuO3 layered structure and obtain higher-stability lithium-rich layered oxides. Using in situ X-ray diffraction and X-ray absorption fine structure, we conclude that oxygen redox was tuned to a reversible level without reduction of Ru as reported previously. The intrinsic mechanism of the modification was further determined by density functional theory calculations. This work will provide a new scope for the strategy of balancing the high capacity and good structural stability of lithium-rich layered oxides.

1. INTRODUCTION The fast development of the electrical vehicle market necessitates high energy storage requirements for lithium-ion batteries. In light of this, high-capacity cathodes are predominantly crucial for the realization of high-energy density batteries. Rock-salt layered oxides LiMO2 (M = Co, Ni, CoxNiyMnz, etc.) are a common kind of cathode with capacities of 140−200 mAh/g1−4 but are not fit for the challenge. Lithium-rich layered manganese-based oxides Li1+xNiyCozMn1−x−y−zO2 are the well-known high-capacity cathodes, but with deficiencies like voltage decay and capacity fading.5−7 Though tremendous efforts have been made to modify the Li1+xNiyCozMn1−x−y−zO2 electrodes,8−12 their practical application is still far from possible. In recent years, many groups have studied new lithium-rich layered oxides in view of their high capacities, such as the Ru-based, Nb-based, and Mo-based lithium-rich oxides.13−17 Tarascon’s group found that the high capacity of lithium-rich layered oxides is realized by successive participation of cationic and anionic redox processes, based on several studies of lithium ruthenium oxides.13,14 These lithium-rich materials all have high capacities exceeding 200 mAh/g and even 300 mAh/g, exhibiting high potential for application in high-energy density lithium-ion batteries. The family of lithium-rich layered materials is now abundant, with various new materials having been proposed,18−22 and opens a new research hot spot for cathode materials. However, among these lithium-rich layered oxides, tremendous efforts should be dedicated to stabilizing their structure during deep charge and discharge. Because of their involvement in anionic redox, deep charge and discharge for high capacity will damage the crystal framework because anions like oxygen ions will be © 2017 American Chemical Society

excessively oxidized. As a result, all of the lithium-rich materials share a common feature in that the cycle stability is very poor,19−21 which strongly attenuates their high-capacity advantage versus conventional cathode electrodes. Therefore, it is crucial to enhance the stability or reversibility of anionic redox in the lithium-rich materials system. Several studies have probed and discussed the intrinsic mechanism involving the reversibility of oxygen redox proposed by the groups of Tarascon and Doublet.21,23 They proposed a metal-driven reductive coupling mechanism to explain the partial reversibility of oxygen redox in Li2RuO3-based oxides in finding that the metal will be reduced finally at the end of the charge process. From their perspectives, the M (3d)−O (2p) (M denotes transition metal ions) covalent bond or the overlap between M (3d) and O (2p) states is responsible for the triggering of the reductive coupling mechanism. This theory is highly consistent with the reduction of Ni ions in lithium-rich manganese-based layered materials as observed previously by our group and other groups.24,25 However, it is hard to explain why Co is not reduced according to this mechanism when large numbers of holes are created on the 2p states of oxygen ions in LiCoO2 under total delithiation, though the covalence of Co− O is sufficiently strong.26−29 However, in our previous work, we have found from the perspective of geometric structure that excess lithium ions in the transition metal layer in lithium-rich oxides will enhance the flexibility and robustness of the local structure, which will facilitate the reversibility of oxygen redox.30 Therefore, the triggering of the reductive coupling Received: November 7, 2016 Revised: February 27, 2017 Published: February 28, 2017 2811

DOI: 10.1021/acs.chemmater.6b04743 Chem. Mater. 2017, 29, 2811−2818

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cut into several small pieces as individual cathode electrodes, followed by drying in a vacuum oven at 120 °C for at least 12 h before the electrodes were fabricated into cells. 2.4. X-ray Absorption Spectroscopy. In situ Ru K-edge XAFS spectra were recorded at beamline 1W1B of the Beijing Synchrotron Radiation Facility (BSRF) and beamline BL14W1 of the Shanghai Synchrotron Radiation Facility using a Si (311) double-crystal monochromator. The in situ cells were fabricated on the basis of the cell used for the in situ XRD measurements; two beryllium windows was placed on both the cathode side and the anode side to allow the Xray beam to penetrate. The spectra was recorded when the cell was charged or discharged to a certain voltage at a slow current density of 30 mA g−1. O K-edge and B K-edge XAS were collected from beamline 4B9B of BSRF and beamline BL10B of National Synchrotron Radiation Laboratory (NSRL). All data processing performed prior to analysis, including energy calibration, background removal, normalization, and Fourier transformation, was performed using Athena. 2.5. DFT+U Calculations. DFT calculations were performed by the Vienna ab initio simulation package (VASP) using the projector augmented-wave (PAW) approach. The spin-polarized generalized gradient approximation (GGA) was employed to investigate the structures, and the Perdew−Burke−Ernzerhof exchange-correlation functional was applied. The energy cutoff for the plane waves was set to 600 eV, and a Monkhorst−Pack scheme 4 × 2 × 2 k-point mesh was used for integration in the irreducible Brillouin zone. The results were obtained within the GGA+U scheme with a Ru Hubbard U value of 5 eV.

mechanism may also be concerned with the excess lithium feature of the lithium-rich layered oxides, which explains its absence in LiCoO2 that has no additional lithium ions in the transition metal layer. However, though the reversibility of oxygen redox can be realized by a reductive coupling mechanism, the excessive oxidation of oxygen still leads to large irreversible capacity. To obtain reversible anionic redox for lithium-rich high-capacity electrodes, tuning the electronic structure between the metal and ligands is an effective way. Here, on the basis of a flexible structure of lithium-rich ruthenium oxide, we found that the reversibility of oxygen redox can be controlled by tuning the electronic structure between Ru and O, which was realized by boron doping in the interstitial sites of the layered crystal structure. The doped samples show a more stable structure and present different charge and discharge behaviors as characterized by in situ X-ray diffraction (XRD) patterns. The Ru Kedge XAFS spectra demonstrate that the reductive coupling mechanism did not happen without the valence of Ru shifting back. In combination with the density functional theory (DFT) calculation, we believe that the introduction of boron atoms changes the electronic structure of the Ru−O system and lowers the energy of the oxygen 2p band. The building of an M−L−X (e.g., Ru−O−B) system that lowers the energy of the local structure is responsible for suppressing the reductive coupling process. This work provides a general strategy for tuning the reversibility of anionic redox in obtaining highcapacity electrodes and opens a new door for the design of electrodes.

3. RESULTS AND DISCUSSION 3.1. General Characterization. All of the samples were synthesized by sintering the mixture of Li2CO3, RuO2, and H3BO3 in a proper proportion as shown in the Experimental Section. To investigate the obtained products, XRD patterns were applied to study the crystal structure as shown in Figure 1a. All of the patterns show the perfect layered feature of the

2. EXPERIMENTAL SECTION 2.1. Sample Preparation. Li2RuO3 and Bx-LRO (x = 0.05, 0.1, and 0.15; for the sake of convenience, the samples doped with different amounts of boron are designated as B0.05-LRO, B0.10-LRO, and B0.15-LRO, corresponding to theoretical B/Ru atomic ratios of 0.05, 0.10, and 0.15, respectively) samples were prepared through a high-temperature solid-state method. Stoichiometric amounts of RuO2·xH2O (J&K, 54% Ru), H3BO3 (J&K, 99.99%), and Li2CO3 (10, 30, 50, and 70% excesses for Li2RuO3, B0.05-LRO, B0.10-LRO, and B0.15-LRO, respectively; Sinopharm, 99.99%) were mixed together and ground homogeneously after the crystal water had been removed at 300 °C for 6 h. The mixture was then transferred to a silica tube furnace and calcined at 950 °C for 12 h and then at 1000 °C for 12 h, and the reaction was naturally quenched at room temperature. 2.2. XRD and X-ray Photoelectron Spectroscopy (XPS) Analysis. The XRD spectra were recorded with a D8-Advance diffractometer (Bruker) with a Cu Kα radiation (λ = 1.5406 Å) source, operated at 40 kV and 40 mA. The spectra were collected in the 2θ range of 10−80°, with a step size of 0.02° and a counting time of 0.5 s. In situ XRD patterns were measured with an in situ stainless steel cell, which was fabricated by using a beryllium window as the current collector to allow the penetration of the X-rays. The cathodes made of Li2RuO3 and B0.15-LRO were just covered by the beryllium window. A glass fiber (GF/D) made by Whatman was used as the separator and lithium metal as the anode. In situ XRD data were collected with the cell being simultaneously charged and discharged at a current rate of 20 mA/g. Rietveld refinement was performed via TOPAS software. XPS data were collected with an AXIS-Ultra instrument from Kratos Analytical (Al Kα radiation; hν = 1486.6 eV). The fitting of the spectra was performed with XPSPEAK41. 2.3. Electrochemical Test. The cathode electrodes were prepared via mixing 80 wt % active materials, 10 wt % polytetrafluoroethylene (PTFE) as the binder, and 10 wt % Super P as the conductive additive. Then the mixture were pasted on aluminum foil, and the foil was rolled into a slice with a thickness of 25−30 μm. Then the slice were

Figure 1. General characterization of Li2RuO3 and Bx-LRO (x = 0.05, 0.1, and 0.15) samples. (a) XRD characterization. (b) Ru K-edge XAFS spectra. (c) O K-edge XAS spectra. (d) B K-edge XAS spectra.

C2/c space group without obvious impurity after the incorporation of certain amounts of boron. However, the (002) peak shows consecutive shifts to smaller angles, as reflected by the inset of Figure 1a. This is probably caused by the doping of boron atoms, which expands the space between the layers. Nevertheless, we did not find any change in the relative intensities of the peaks, suggesting that boron doping 2812

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refinement analysis of the XRD patterns was performed to investigate the doping site. By utilizing the Li2RuO3 model, the refinement converged to a result with an Rwp of 6.49% (Figure 2a). After the boron atoms had been introduced into the

did not change the long-range order. Because of the facilitation of boron in the calcination, the whole intensity of the diffraction peaks increases with the increase of the amounts of boron, indicating the degree of crystallinity is higher after boron doping. The samples with boron contents higher than that of B0.15-LRO were also prepared, however, with an impurity like RuO2 present. Therefore, the subsequent electronic and structural analyses of these impure samples are excluded. The electronic structural variation of Li2RuO3 after the incorporation of boron atoms was studied. The variation of the electronic structure of Ru is characterized by Ru K-edge XAFS, as shown in Figure 1b. The absorption edge of Ru has a tendency to shift to a lower energy as the boron content increases, presenting a reasonable evolution as electric neutrality requires a lower valence of Ru to compensate for the charge brought by boron atoms. To investigate the electronic structural change of oxygen ions, we performed the soft X-ray absorption spectroscopy (XAS) technique in obtaining the O K-edge XAS spectra, as shown in Figure 1c. The pre-edge peak of O K-edge XAS spectra corresponds to the transition of O 1s electrons to Ru (4d)−O (2p) hybridized states.29,31 As shown in Figure 1c, the shape and intensity of the O K-edge spectra vary greatly with the amount of boron doping. The shoulder peak at ∼535 eV for the pristine sample attenuates for B0.05-LRO and then disappears for B0.1-LRO and B0.15-LRO. Because of the screening effect,28 the disappearance of the shoulder peak at higher energies indicates the increase in the number of electrons on the O 2p band or the decrease in the number of the holes of the O 2p band, which suggests the decrease in the covalence of the Ru−O bond due to the incorporation of boron atoms. From B0.1LRO to B0.15-LRO, the area of the normalized pre-edge peak decreases (see the inset of Figure 1c), showing a further decrease in the number of electron holes on the O 2p band, as does the covalence of the Ru−O bond.31 All of these electronic structural changes provide potent evidence of the bulk doping of the boron atoms after we excluded the possibility of Li or Ru deficiency in compensating for the charge of boron (see sections I and II of the Supporting Information). It is also very crucial to determine the actual positions of doped boron atoms in the structure. The most probable case is that in which boron atoms reside in the pseudotetrahedral interstitial sites of Li2RuO3. Boron atoms are generally considered to be coordinated with three or four oxygen ions, and no evidence of a BO6 octahedron has been found, which suggests boron atoms are not likely to substitute for Ru ions. Because the radius of B3+ ions is only 0.027 nm, interstitial-site doping is possible, as revealed by previous works that examined B-doped TiO2.32,33 Previous investigation of boron-doped LiCoO234 and Li-rich manganese-based materials25 further proves that the doping of B in tetrahedral interstitial sites of a layered rock-salt type structure is possible. To probe the actual state of the doped boron atoms, B K-edge XAS spectra were recorded as shown in Figure 1d. The signal of boron become stronger as the boron content increased. The sharp peak A at ∼194 eV corresponds to the trigonal BO3 group. The broader peak B at 197−199 eV corresponds to the BO4 group. The broad peak C is contributed by the BO4 group at 200−202 eV and the BO3 group at 203−204 eV jointly, as inferred from previous work.35−37 We can determine that in the layered structure boron atoms formed as both BO3 and BO4 groups and stayed as local polyanions. Furthermore, the Rietveld

Figure 2. Rietveld refinement of the XRD patterns of the B0.15-LRO sample. (a) Refinement with the model of Li2RuO3 without any boron. (b) Refinement with the model of Li2RuB0.15O3 with boron atoms in the pseudotetrahedral interstitial site.

interstitial site, the refinement was improved with an Rwp of 6.06% (Figure 2b). The refined atomic information on the supercell of Li16Ru8BO24 is shown in Table S2. Although boron atoms have a small scattering cross section and influence the XRD reflections to a limited extent, the improvement in the Rwp value still positively supports the existence of boron atoms in the interstitial sites of layered Li2RuO3. 3.2. Electrochemical Performance. The electrochemical performance of Li2RuO3 and its boron-doped samples was tested in galvanostatic charge−discharge mode. The first and second charge and discharge curves and derived dQ/dV plots are presented in Figure 3. The first charge process of Li2RuO3 has previously been shown to consist of two processes involving Ru redox and oxygen redox, corresponding to the 3.6 and 4.2 V voltage platforms, respectively,25 as shown in Figure 3a. However, after the doping of boron, the first charge behavior has been largely changed so that the 3.6 V platform stretches while the platform of oxygen redox is shortened, as observed from the charge curves of Bx-LRO (x = 0.05, 0.1, and 0.15) in Figure 3a. This suggests that Ru takes more responsibility for the charge compensation process while oxygen takes less, though the total capacity decreases with the increase in boron content. Other than that, the voltage of the oxygen redox plateau increased from 4.2 to 4.3−4.4 V after boron doping, as shown by the dQ/dV plots in Figure 3a. This enhancement is consistent with the inductive effect of polyanions proposed by Goodenough in polyanion type cathodes like LiFePO4.38 It should be noted that though the plateaus are lifted, the redox processes of Ru and oxygen are totally completed before 4.6 V, not disrupted. This is crucial for the comparison between Li2RuO3 and Bx-LRO samples. What interests us more is the fact that voltage plateaus in the first charge process are straighter than those in the pristine one, as reflected by the sharper peaks of the reduction and oxidation 2813

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more perfect with an increase in boron content than that corresponding to peak B at lower voltages. More than that, the oxygen redox plateau at ∼4.3 V is still observed for Bx-LRO samples in the second cycle, demonstrating structural retention and redox reversibility that are better than those of pristine Li2RuO3. We also provide the charge/ discharge performance at the 15th cycle (Figure S2), which demonstrates a similar difference between the pristine and BxLRO samples. It is obvious that the electronic structural variation is responsible for the difference between the samples. With additional boron, the capacity is lower but the stability higher, as proven by the cycling performance shown in Figure 4. For boron-doped samples, the capacity of the first cycle is

Figure 4. Cycling performance of the electrodes.

lower than that of second cycle, which was probably caused by the activation effect due to the particles being larger than the pristine sample. Though a smaller particle size may enhance the cycling stability in most cases, such as the Si anode, in our work, boron-doped samples with larger particles demonstrate a cycling ability that is better than that of pristine Li2RuO3. This further emphasizes the effect of boron doping in enhancing cycling stability. We speculate that the first charge process of Li2RuO3 will lead to the total rearrangement of the structure after the loss of oxygen while being inhibited by the tuning of the electronic structure with boron atoms. Only in such a situation can the structure of boron-doped layered oxides be reserved without any oxygen loss. 3.3. In Situ XRD Characterization. To verify our speculation, we used in situ XRD to track the long-range order evolution of the structure before and after boron doping. Figure 5 shows the in situ XRD results of Li2RuO3 and B0.15LRO between 17.5° and 20° for the first cycle and second charge process, respectively. The first charge process of Li2RuO3 shows a continuous three-phase transition feature of the (002) peak in the first plateau as shown in Figure 5a, with three interphases labeled P1−P3 between Li2RuO3 and RuO3. However, for B0.15-LRO, the phase transition process demonstrates behavior quite distinctive compared to that of the P1 phase, which was a longer process, and the P2 phase appeared to dependent on the process of the P3 phase transition rather than an independent phase change presented in Li2RuO3, as shown in Figure 5b. This difference may be induced by the variation in the electronic structure caused by the incorporation of boron. During the second plateau of the first charge process, a solid solution structural evolution was presented for both Li2RuO3 and B0.15-LRO. However, the (002) peak of Li2RuO3 shifts by 0.25°, while B0.15-LRO shows little change, suggesting the latter undergoes less structural destruction as a result of the oxygen redox process. Other than

Figure 3. Electrochemical performance of Li2RuO3 and Bx-LRO (x = 0.05, 0.1, and 0.15). (a) First charge and discharge curves and dQ/dV plots. (b) Second charge and discharge curves and dQ/dV plots.

in dQ/dV plots. For the first discharge curves, the pristine sample shows a slope feature with a volcano type peak displayed in the dQ/dV plot, while the doped samples present two independent plateaus and two distinctive needlelike reductive peaks in the dQ/dV plots. These differences in charge/discharge behavior are more evident in the second charge and discharge processes, as shown in Figure 3b. The charge and discharge curves both exhibit distinctive voltage plateaus and oxidative and reductive peaks for the boron-doped samples, compared with that of pristine Li2RuO3. All of these changes are supposed to be caused by boron doping, which induces the variation of the electronic structure and thus the electrochemical behavior. The relative size of the two peaks (labeled A and B) in the dQ/dV plots during the discharge of the second cycle shown in Figure 3b changes as a function of boron content, especially for the B0.15-LRO sample; specifically, the intensity of peak A increases relative to that of peak B as the boron content increases. Because the peak intensity of the dQ/dV plot is in direct relation to the slope of the voltage platform, the relative change in the size of the dQ/ dV plots suggests variation of the voltage platform, which might result from the influence of the two-phase transition process caused by doping different amounts of boron. Specifically, the two-phase transition reaction corresponding to peak A becomes 2814

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results of in situ XRD are consistent with our previous speculation about the electrochemical performance and let us conclude that boron doping tuned the reversibility of oxygen redox and thus the structure. 3.4. In Situ Ru K-Edge XAFS. To further explore the redox process of the total charge and discharge process before and after doping, we performed the in situ Ru K-edge XAFS measurement, shown in Figure 6. During the first charge

Figure 5. In situ XRD results for Li2RuO3 and B0.15-LRO during the first cycle and second charge process. (a) In situ XRD of the first charge process of Li2RuO3. (b) In situ XRD of the first charge process of B0.15-LRO. (c) In situ XRD of the first discharge process of Li2RuO3. (d) In situ XRD of the first discharge process of B0.15-LRO. (e) In situ XRD of the second charge process of Li2RuO3. (f) In situ XRD of the second charge process of B0.15-LRO.

that, the broadening of the (002) peak of Li 2 RuO 3 demonstrates that a large internal stress existed in the crystal39 due to the induced structural distortion in the oxygen redox process. There is no obvious broadening of the (002) peak in the B0.15-LRO sample (Figure 5b), demonstrating little internal stress and the higher stability of the modified layered structure. Panels c and d of Figure 5 present the in situ XRD results of the first discharge process of Li2RuO3 and B0.15-LRO, respectively. The evolution of the (002) peak of Li2RuO3 presents a two-stage feature: before 3.3 V, the (002) peak continuously substantially shifts during the oxygen redox process, demonstrating a solid solution behavior; after that, no shifting occurs, but the peak becomes broadened. However, for B0.15-LRO, peak evolution shows a two-phase transition behavior rather than a solid solution change before 3.3 V. The difference in structural evolution between the first charge process and the discharge process of Li2RuO3 is caused by the structural reorganization caused by the loss of oxygen, resembling the situation seen for Li-rich manganese-based materials.40 Here, the B0.15-LRO remaining as a two-phase transition change in the discharge process can be the route for better maintaining the layered structure during the oxygen redox process or reflect the better reversibility of the oxygen redox compared with Li2RuO3. Such a difference also extends to the second charge process (see Figure 5e,f). By comparing the initial position of the (002) peak with that at the end of the charge and discharge processes, we also can find that the reversibility of B0.15-LRO is better than thta of LRO. These

Figure 6. In situ Ru K-edge XAFS spectra results. (a) In situ Ru K-edge XAFS spectra of Li2RuO3 during the first charge process. (b) Fouriertransformed EXAFS results of Li2RuO3 for the first charge process. (c) In situ Ru K-edge XAFS spectra of B0.15-LRO during the first charge process. (d) Fourier-transformed EXAFS results of B0.15-LRO for the first charge process. (e) In situ Ru K-edge XAFS spectra of Li2RuO3 during the first discharge process. (f) Fourier-transformed EXAFS results of Li2RuO3 for the first discharge process. (g) In situ Ru K-edge XAFS spectra of B0.15-LRO during the first discharge process. (h) Fourier-transformed EXAFS results of B0.15-LRO for the first discharge process.

process, the absorption edge of the Ru K-edge shifts to a higher energy before 4.2 V but shifts to a lower energy when it was charged to 4.6 V (Figure 6a; the edge position change as a function of capacity is shown in Figure S3a). This result is consistent with the previous work and was determined to be the reductive coupling mechanism by Tarascon’s group.13,23 Correspondingly, the intensity of the peaks of Fouriertransformed EXAFS spectrum, denoting the local structural disorder, shows similar evolution with electronic structure 2815

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Chemistry of Materials (Figure 6b). The two peaks of the Ru−O shell and Ru−Ru shell all demonstrate the increased intensity before 4.2 V, while after that, the intensity decreases drastically, indicating the disordered local structure was induced by the reductive coupling mechanism (Figure S3b). However, for the first charge process of B0.15-LRO, the Ru K-edge continuously shifts to a higher energy without any fluctuation (Figure S3a), which is different from the case for Li2RuO3. Analogously, the evolution of the radial distribution functions (Figure 6d and Figure S3b) is generally monotonous over the whole process with little exception, suggesting that the local structure is becoming ordered during the total charge process, even within the oxygen redox process. The electronic structural change in oxygen for B0.15-LRO was characterized by O 1s XPS analysis (Figure S4), which exhibits a reversible oxygen redox during the charge and discharge processes. Such a difference between Li2RuO3 and B0.15-LRO is also determined in the discharge process, as seen in Figure 6e−h. The valence of Ru ions in Li2RuO3 increases first upon discharge from 4.6 to 3.5 V and then decreases after that, while a monotonous increase is seen for B0.15-LRO. The intensities of the peaks in the radial distribution functions of Li2RuO3 shown in Figure 6f tend to increase first and then decrease, while for B0.15-LRO, the total process exhibits a monotonous increase. This further demonstrates the absence of a reductive coupling mechanism in B0.15-LRO though in the presence of oxygen redox. As we have investigated previously,30 the whole charge process of Li2RuO3 can be divided into two stages. In the lower voltage range called stage L, Ru ions are oxidized with the structure being more ordered; in the higher voltage range called stage H, Ru ions are reduced successively and O ions are oxidized, with the structure being more disordered. In stage H, the reductive coupling mechanism is used when the oxygens are oxidized to a certain extent, and the local structure will be distorted to relieve the energy. However, after the doping of boron, the reductive coupling mechanism was replaced with a monotonous electronic structural change, with the local structure being more ordered during the charge process. Because the reduction of Ru ions begins at the start of the oxidation of oxygen, the variation of the charge compensation mechanism should be contributed by the electronic structural change after doping, though the oxygen redox process was limited to a smaller extent. The modified Ru−O−B structure can make itself stable under oxygen redox without a reductive coupling mechanism; thus, Ru ion are oxidized successively, and the structure tends to be ordered without any distortion. 3.5. DFT Calculation. To further investigate the intrinsic mechanism leading to the change in the redox process, we performed DFT calculations to theoretically answer such a question. The calculation was realized with a supercell of Li16Ru8O24 and Li16Ru8BO24, in which the boron atom resides in the pseudotetrahedral interstitial site, and the optimized structures are shown in panels a and c of Figure 7, respectively. After total delithiation, the optimized structure of Ru8O24 demonstrates a large distortion as shown in Figure 7b, with O−O dimers formed as the O22−-like species and the O−O bond being 2.418 Å long. The local view of optimized Ru8BO24 (the optimization details are shown in section III of the Supporting Information) is shown in Figure 7d, with less distortion compared with Ru8O24, and the average length of O−O bonds is ∼2.515 Å, which is longer than that of Ru8O24. This implies that the oxidation of oxygen in B-doped Li2RuO3

Figure 7. DFT calculation results. (a) Schematic of the optimized crystal model of Li16Ru8O24. (b) Local structure of the calculated Ru8O24 viewed from the c axis. (c) Schematic of the optimized crystal model of Li16Ru8BO24. (d) Local structure of the calculated Ru8O24 viewed from the c axis. (e) Calculated PDOS of Li16Ru8O24. (f) Calculated PDOS of Li16Ru8BO24.

is limited, and the structure exhibits a higher order, which confirms the experimental results. To probe the inherent mechanism, we plot the partial density of states (PDOS) to observe the electronic structural change before and after boron doping, as shown in panels e and f of Figure 7. The majority of the electrons that reside around the Fermi level belong to Ru, and oxygen overlaps with Ru between the energy gap of (−2, 0). Sequentially, the Ru ion will first distribute its electrons upon delithiation, and then oxygen, via the reductive coupling mechanism. After boron doping, the occupation of oxygen electrons between the energy gap of (−2, 0) was reduced from 50.12 to 46.32%, suggesting the extent of overlap between the Ru 4d and O 2p states is reduced. This result is consistent with the previous electronic structural characterization that revealed that the covalence of Ru−O is reduced. As a result, Ru ions will take more responsibility for the redox reaction, and oxygen redox will be limited to a certain extent. To determine why the reductive coupling mechanism is absent in boron-doped samples, we compared the PDOS before and after structural optimization [demonstrated as rigid and relaxed structure, respectively (see Figure S6)], as mentioned in ref 23. The results show that for half-delithiated LiRuO3, the relaxed density of states (DOS) demonstrates no obvious change compared with the rigid structure DOS, suggesting LiRuO3 is a self-stable structure. For B-LiRuO3, the rigid structure and relaxed DOS all present an energy of the oxygen 2p states 2816

DOI: 10.1021/acs.chemmater.6b04743 Chem. Mater. 2017, 29, 2811−2818

Chemistry of Materials comparative lower than that of LiRuO3, indicating the higher stability of the oxygen. With respect to totally delithiated RuO3, the relaxed DOS demonstrates a large variation compared with the rigid structure DOS whereby the system changes from a conductor to a semiconductor with an energy gap (Figure S7). Such electronic structural rearrangement is realized by the reductive coupling mechanism leading to structural distortion and to the decreased energy of the oxygen 2p electrons, thus maintaining the whole layered structure. However, the borondoped sample after total delithiation (B-RuO3) just shows an energy decrease for the comparison of rigid structure and relaxed DOS, still remaining a conductor without electronic structural rearrangement. This suggests that the B-RuO3 is a stable structure rather than a metastable structure after a certain degree of oxidation of oxygen due to the incorporation of boron.

ACKNOWLEDGMENTS

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b04743. XPS semiquantitative analysis of B0.15-LRO (section I), XRD refinement (section II), and DFT calculation (section III) (PDF)

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The authors appreciate the help with the difficult XAS measurement from the staff of beamline 1W1B of the Beijing Synchrotron Radiation Facility (BSRF) and beamline BL14W1 of the Shanghai Synchrotron Radiation Facility (SSRF), as well as the help with the soft XAS measurement from the staff of beamline 4B9B of BSRF and beamline BL10B of the National Synchrotron Radiation Laboratory (NSRL). This work was financially supported by the new energy project for electric vehicle of national key research and development program (2016YFB0100200) and the National Natural Science Foundation of China (51671004).

4. CONCLUSIONS The stability of the layered structure is related to the reversibility of oxygen redox when the lithium-rich layered oxides deliver a high capacity, though with the reductive coupling mechanism, the excessive oxidation of the oxygen leads to a significant decrease in capacity during cycles. After boron doping, oxygen redox was tuned to the proper extent with the totally delithiated structure stable, without being distorted to relieve the energy. Then the materials show better reversibility of structural evolution during the charge and discharge process and deliver excellent cycling stability. The central goal of this work is to provide a general strategy for balancing the capacity and stability by tuning the reversibility of anionic redox in lithium-rich high-capacity electrodes. This balance means an acceptable capacity as well as its stability during cycles. Simply decreasing the cutoff voltage is an alternative strategy but to some extent may not be as good as tuning its intrinsic electronic structure. In general, through this work, we demonstrate that oxygen redox can be tuned by modifying the electronic structure of the metal−oxygen group and thus realize the balance between high capacity and high stability. For most of the lithium-rich layered oxides, the high capacities cannot be sustained during long cycles because of the excessive oxidation of oxygen ions. This work will provide a route for the modification of these lithiumrich layered systems to obtain stable high-capacity electrodes.

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Article

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Dingguo Xia: 0000-0003-2191-236X Notes

The authors declare no competing financial interest. 2817

DOI: 10.1021/acs.chemmater.6b04743 Chem. Mater. 2017, 29, 2811−2818

Article

Chemistry of Materials

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DOI: 10.1021/acs.chemmater.6b04743 Chem. Mater. 2017, 29, 2811−2818